design of high-strength bainitic...

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Design of novel high-strength bainitic steels. Part II

F. G. Caballero, H. K. D. H. Bhadeshia, K. J. A. Mawella, D. G. Jones and P. Brown

Dr F. G. Caballero and Professor H. K. D. H. Bhadeshia are in the Department of

Materials Science and Metallurgy, University of Cambridge, Pembroke Street,

Cambridge CB2 3QZ, UK. Dr K. J. A. Mawella, D. G. Jones and Dr P. Brown are in

the Structural Materials Centre, Defence Evaluation Research Agency, R1079,

Bldg A7, Farnborough GU14 0LX, UK

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A combination of thermodynamic, kinetic and mechanical property models and

physical metallurgy principles was used in Part I of this study, to propose a number

of alloys which exploit the carbide-free bainitic microstructure at its theoretical best.

These alloys have been manufactured and the present paper (Part II) reports the

results of metallographic characterisation and mechanical tests. The proposed steels

are found to have the highest ever combination of strength and toughness for bainitic

microstructures, matching even the maraging steels which are at least thirty times

more expensive. The work confirms the alloy design procedures explained in Part I.

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1 Introduction

A number of carbide-free bainitic steels were proposed for manufacturing in Part I of

this study. The steels were designed in the basis that the fraction of bainitic ferrite

obtained by continuous cooling transformation should be high enough to avoid any

large and unstable regions of high-carbon retained austenite, having only the films of

retained austenite to separate the bainite platelets. The steels also had to meet certain

hardenability and strength requirements.

The work presented here (Part II) is concerned with the microstructural and

mechanical property characterisation of the three bainitic steels proposed in Part I of

this study. As will be seen later, the results are exciting.

2 Experimental procedure

The actual chemical compositions of the steels studied are given in Table 1. Alloys

were prepared as 35 kg vacuum induction melts using high purity base materials.

After casting and cropping, the ingots were hot forged down to a thickness of 65 mm.

These were then homogenised at 1200 oC for 2 days and cut into smaller 65 mm

thick samples which were then forged down to 50 mm thickness. These samples

were then held at 900 oC for 2 hours, removed from the furnace and immediately hot-

pressed to a thickness of 25 mm before their temperature fell below 750 oC. Finally

they were allowed to cool in air. The particular deformation route used here (Fig. 1)

is consistent with a proprietary manufacturing standard.

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2.1 MICROSTRUCTURAL EXAMINATION OF STEELS

Quantitative X-ray analysis was used to determined the fraction of retained austenite.

For this purpose, samples were cut from undeformed regions of Charpy specimens.

After grinding and final polishing using 0.25 µm diamond paste, the samples were

etched to obtain an undeformed surface. They were then step-scanned in a Philips -

PW 1730 X-ray diffractometer using unfiltered Cu Kα radiation. The scanning speed

(2θ) was 1 degree min-1. The machine was operated at 40 kV and 40 mA. The

retained austenite content was calculated from the integrated intensities of (200),

(220) and (311) austenite peaks, and those of (002), (112) and (022) planes of

ferrite.1 Using three peaks from each phase avoids biasing the results due to any

crystallographic texture in the samples.2 The carbon concentration in the austenite

was estimated by using the lattice parameters of the retained austenite.3

Specimens for transmission electron microscopy (TEM) were machined down to

3 mm diameter rod. The rods were sliced into 100 µm thick discs and subsequently

ground down to foils of 50 µm thickness on wet 800 grit silicon carbide paper. These

foils were finally electropolished at room temperature until perforation occurred,

using a twin-jet electropolisher set at a voltage of 40 V. The electrolyte consisted of

5 % perchloric acid, 15 % glycerol and 80 % methanol. The foils were examined in a

JEOL JEM-200 CX transmission electron microscope at an operating voltage of

200 kV.

Optical and scanning electron microscopy (SEM) were used to examine the etched

microstructures. Specimens were polished in the usual way and etched in 2% Nital

solution, and examined using a JEOL JXA-820 scanning electron microscope

operated at 10-15 kV. The volume fraction of bainite (Vb) was estimated by a

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systematic manual point-counting procedure on scanning electron micrographs.4 A

grid superimposed on the microstructure provides, after a suitable number of

placements, an unbiased statistical estimate of Vb.

2.2 MECHANICAL PROPERTIES DETERMINATION

Tensile testing was carried out in accordance with BS EN 10 002-1: 1990 at room

temperature on a 100 kN Instron-6025 machine at a crosshead speed of 2 mm min-1.

Two specimens were tested for each alloy.

Impact toughness was measured at temperatures between 20 and -120 oC using a

300 J Charpy testing machine. Specimens were tested in accordance with BS EN 10

045-1: 1990. Six specimens were tested at each temperature for every alloy.

Compact tension specimens were used to measure values of plane strain fracture

toughness (KIc) at room temperature for the Ni1 and Ni2 alloys in accordance with

BS 7448 Part1: 1991. Compact tension / fracture toughness testing was done on a

100 kN ESH servohydraulic machine. The crosshead speed used was 1 mm min-1.

Fractography was carried on the Charpy impact toughness specimens using a JEOL

JXA-820 scanning electron microscope operating at 20 kV.

2.3 TEMPERED MICROSTRUCTURES

Austenite formation begins during heating at the Ac1 temperature and is completed

when the Ac3 temperature is reached. Tempering must be carried out below the Ac1

temperature to avoid the accidental formation of austenite. The austenite formation

temperatures were therefore determined using a Thermecmastor-Z thermomechanical

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simulator. Cylindrical specimens 12 mm in height and 8 mm in diameter were heated

at a rate of 10 oC s-1 to 1000 oC and then cooled at 10 oC s-1. The Formation of

austenite during heating was detected by monitoring the fractional change in

dilatation with temperature. Table 2 shows the Ac1 and Ac3 temperatures. Note that

the Ac1 temperature is ill-defined in the present context because the microstructures

already contain some retained austenite. Ac1 should therefore be interpreted to mean

the temperature at which austenite growth begins during heating at the specified rate.

Specimens of the three alloys were tempered at temperatures ranging from 400 to

700 oC for an hour. Specimens for TEM were produced as described above.

3 Results and discussion

3.1 CHARACTERISATION OF MICROSTRUCTURE

Experimental data of the designed steel microstructures are presented in Table 3.

These data and the micrographs presented in Fig. 2 reveal that Ni1 and Ni2 have the

desired microstructure consisting of mainly bainitic ferrite and retained austenite.

Because of the high carbon content in austenite (xγ) for these alloys (Table 3), the

majority of residual austenite present after bainite is formed, is retained on cooling to

room temperature. In the SEM micrographs martensite regions are relatively

unetched, and appear light in colour. The thin interlath films of austenite, however,

have been etched away. It is difficult in thin-foil TEM experiments to judge whether

any martensite has formed as a result of the thinning process or whether it was

present in the original microstructure. To cope with this, the samples were tempered

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at 200 oC for 2 h before making foils. Thus, martensite present in the original sample

would contain carbides, whereas any untempered martensite can be interpreted to be

austenite which underwent transformation during the foil preparation process. The

observation of tempered microstructures (Fig. 2.c) showed that the unetched regions

in Fig 2.b are martensitic. Note that the low temperature tempering (200 oC) heat

treatment does not affect bainitic ferrite since the latter does not contain excess

carbon.

Due to the high volume fraction of bainitic ferrite in Ni1 and Ni2 alloys, 0.62 and

0.81 respectively, the retained austenite was largely present as films between the

subunits of bainitic ferrite. Figure 3 shows typical bright-field images of carbide-free

upper bainite in Ni1 and Ni2 with interlath retained austenite films. These films have

a typical wavy morphology characteristic of the bainite in high-silicon steels

(Fig. 3.b).5-8

By contrast, the small amount of bainite in the Mn alloy (Table 3) causes much of the

residual austenite to transform to martensite during cooling because of lower carbon-

enrichment of the austenite (0.55 wt-%, Table 3). There is only 10 % of the residual

austenite retained to room temperature (Vγ=0.07, Table 3). These results are

consistent with the data on instability of residual austenite as function of austenite

carbon content illustrated elsewhere.9

The results of hardness tests are presented in Table 4. There is no doubt that the

microstructures of Ni1 and Ni2 alloys in the as-received condition consist mainly of

bainite and retained austenite. Their hardness values are closer to those of

microstructures obtained isothermally at 375 oC than the respective hardness values

of the martensite as determined by water quenched following austenitisation at

1000 oC for 15 min. Any difference between the hardness values of as-received and

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isothermally transformed microstructures in Ni1 and Ni2 alloys is due to the different

degree of transformation to bainitic ferrite. Not surprisingly, the hardness value for

the as-received microstructure of Mn alloy (Fig. 2.a) is similar to that of a fully

martensitic sample.

3.2 MECHANICAL PROPERTIES

3.2.1 TENSILE STRENGTH AND DUCTILITY

Tensile test results are presented in Table 5. Plates of bainitic ferrite are typically

10 µm in length and about 0.2 µm in thickness (Fig. 3). This gives a rather small

mean free path for dislocation glide. Thus, the main microstructural contribution to

the strength of bainite is from the extremely fine grain size of bainitic ferrite.10

It is difficult to separate the effect of retained austenite on strength in these steels

from other factors. Qualitatively, austenite can affect the strength in several ways.

Residual austenite can transform to martensite during cooling to room temperature,

thus increasing the strength as observed in Mn alloy (Tables 3 and 5). On the other

hand, retained austenite interlath films can increase the strength by transforming to

martensite during testing, similar to the behaviour of TRIP (transformation induced

plasticity) steels.

The low yield/ultimate tensile strength ratios (YS/UTS) in Table 5 are due to the

presence of austenite and the generally large dislocation density in the

microstructure.11 Consequently, retained austenite increases the strain-hardening rate

of the steel. Likewise, tensile elongation is controlled by the volume fraction of

retained austenite.12 Retained austenite is a ductile phase compared to the bainitic

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ferrite and would be expected to enhance ductility as far as the austenite is

homogeneously distributed along plate boundaries (film austenite). However,

isolated pools of austenite (blocky austenite) would influence unfavourably on both

elongation and UTS. From Table 5, it is clear that the steels present a combination of

high strength and good ductility. Moreover, tensile properties are much higher than

those required for defence applications (Table 3 in Part I).

3.2.2 IMPACT TOUGHNESS

Charpy impact test results are listed in Table 6 for all the alloys. The levels required

for defence applications in impact toughness have been also well achieved (Table 3

in Part I). A considerable improvement in toughness is obtained when the volume

fraction of bainite increases in the microstructure. This improvement occurs despite

the fact that the strength of the microstructures involved remains almost unchanged

(Table 5). From the data in Table 3, it is evident that the volume fraction of bainite

Vb and thus the carbon content of the retained austenite xγ explains the improvement

in toughness observed in Ni1 and Ni2 alloys as compared with the Mn alloy. The

results are consistent with the enhancement of toughness expected when the amount

of blocky austenite and martensite are reduced and, in general, when the stability of

residual austenite is increased. Figure 4 shows the fracture surfaces of the Charpy

impact specimens tested at room temperature. The two Ni alloys show ductile

dimpled fracture surfaces whereas the Mn alloy, which contains a large amount of

martensite, has many of the features of quasi-cleavage with isolated patches of

ductile fracture.

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3.2.3 FRACTURE TOUGHNESS

A compact tension specimen design was used with thickness B=23.1 mm, width

W=46.5 mm, and crack length, a=24.5 mm, to width ratio of a/W=0.5. With the

measured yield stress of 1100 MPa, this would give a plane strain, KIc, measurement

capacity of ~ 300 MPa m . In the event, none of the specimens failed at sufficiently

low load to satisfy the KIc validity requirement.

The specimens showed a fairly significant non-linearity before maximum load. This

disqualified the use of KQ as a KIc measurement. The values quoted in Table 5 are for

the stress intensity at maximum load, Kmax. All the tests failed by microvoid

coalescence with no sign of cleavage. Tearing was stable at maximum load under

displacement control. Taking account of the non-linearity in the trace up to

maximum load gives the J-Integral at maximum load, Jmax and the stress intensity

KJmax calculated from Jmax values listed in Table 5.

Although the obtained fracture toughness results can not be considered valid, the

results are promising in that the steels tested are so tough that larger samples would

be needed to measure KIc. Toughness values of nearly 130 MPa m have been

obtained for strength in the range of 1600-1700 MPa. The good fracture toughness

obtained in these alloys is attributed to the presence of thin films of thermally and

mechanically stable interlath retained austenite. The role of retained austenite is to

refine the effective fracture grain size and to a blunt propagating crack.13

Figure 5 shows properties of mixed microstructures of bainitic ferrite and austenite,

versus those of quenched and tempered (QT) low-alloy martensitic alloys and

maraging steels.7 Small points in the graph refer to previous work,5-7,9,13 whereas the

two large points correspond to experimental data shown in Table 5. The alloys

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designed theoretically in this work and produced by commercial continuous cooling

present the highest strength/toughness combinations ever recorded in bainitic steels.

These alloys show better mechanical properties of the QT low-alloy martensitic

alloys and match the critical properties of maraging steels, which are at least thirty

times more expensive.

3.3 CHARACTERISATION OF TEMPERED MICROSTRUCTURES

Figure 6 shows a plot of hardness values as a function of tempering temperature for

the three studied alloys. Tempering the designed microstructures of Ni1 and Ni2

steels at temperatures lower than 550 oC for an hour did not result in any significant

loss of hardness. However, tempering at 450 oC for an hour the as-received

microstructure of Mn alloy led to a drop in the hardness value from 597 to 555 HV.

Since the Mn alloy contains predominantly martensite (Vα' = 0.67, Table 3)

tempering at 400 oC for an hour is expected to lead to the precipitation of any excess

carbon. This is confirmed by Fig. 7.a which shows discrete carbide particles

precipitated inside a martensite plate in the Mn alloy. The retained austenite was still

intact with no significant signs of recovery in the bainitic ferrite.

Figures 7.b and c show bright-field images of microstructures obtained by tempering

at 400 oC for an hour in Ni1 and Ni2 alloys. Comparing tempering microstructures

with those as-received conditions (Fig. 3), it is clear that the tempering at 400 oC did

not result in any recovery in the bainitic ferrite (Fig. 7.b) or in the diffusional

decomposition of high carbon retained austenite (Fig. 7.c and d).

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4 Conclusions

It has been demonstrated experimentally that models based on phase transformation

theory can be applied successfully to the design of high strength and tough steels.

These alloys designed to ensure that the hardenability of the steel is sufficient for

industrial production have achieved the highest strength and toughness combinations

to date for bainitic steels, at a cost thirty times less than that of maraging steels.

Likewise, it has been found that these two new steels also show an important

resistant to tempering. Steels have been tempered at temperatures ranging from 400

to 700 oC for an hour and no significant changes in hardness were detected in the

microstructure at tempering temperatures lower than 550 oC.

5 Acknowledgments

This work was carried out as part of Technology Group 4 (Materials and Structures)

of the MoD Corporate Research Programme. The authors would like to thank to

Professor Alan Windle for the provision of laboratory facilities at the University of

Cambridge.

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6 References

1. J. DURNIN and K. A. RIDAL: Journal of the Iron and Steel Institute, 1968, 206,

60.

2. M. J. DICKSON: J. Appl. Cryst., 1969, 2, 176-180.

3. D. J. DYSON and B. HOLMES: Journal of the Iron and Steel Institute, 1970,

208, 469.

4. G. F. VANDER VOORT: ‘Metallography. Principles and Practice’, 427; 1984,

New York, McGraw-Hill.

5. V. T. T. MIIHKINEN and D. V. EDMONDS: Mater. Sci. Technol., 1987, 3, 422-

431.


440.


449.

8. L. C. CHANG: Metall. Trans., 1999, 30A, 909-916.

9. H. K. D. H. BHADESHIA and D. V. EDMONDS: Metal Sci., 1983, 17, , 411-

419.

10. K. J. IRVINE, F. B. PICKERING, W. C. HESELWOOD and M. J. ATKINS: J.

iron Steel Inst., 1957, 195, 54-67.

11. A. P. COLDREN, R. L. CRYDERMAN and M. SEMCHYSHEN: ‘Steel

Strengthening Mechanisms’, 17; 1969, Ann Arbor, USA, Climax Molybdenum.

12. B. P. J. SANDVIK and H. P. NEVALAINEN: Met. Tech., 1981, 15, 213-220.

13. H. K. D. H. BHADESHIA and D. V. EDMONDS: Metal Sci., 1983, 17, 420-425.

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Table 1 Actual Chemical Composition of Designed Alloys, wt-%

Alloy C Si Mn Ni Cr Mo V

Mn 0.32 1.45 1.97 <0.02 1.26 0.26 0.10

Ni1 0.31 1.51 <0.01 3.52 1.44 0.25 0.10

Ni2 0.30 1.51 <0.01 3.53 1.42 0.25 <0.005

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Table 2 Ac1 and Ac3 Temperatures in oC

Alloy Ac1 Ac3

Mn 807 881

Ni1 759 818

Ni2 775 835

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Table 3 Quantitative Data on Microstructure and Hardness*

Alloy Vb Vγ Vα' xγ, wt-% Hardness, HV30

Mn 0.26±0.01 0.07±0.01 0.67±0.02 0.55 597±2

Ni1 0.62±0.05 0.12±0.01 0.26±0.04 0.92 493±5

Ni2 0.81±0.06 0.11±0.01 0.08±0.05 1.03 536±6

Vb bainitic ferrite volume fraction; Vγ retained austenite volume fraction; Vα'

martensite volume fraction; xγ carbon content in austenite

*Hardness values are averaged over 10 tests

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Table 4 Hardness Data of Different Microstructures of the Alloys Studied

Alloy Heat treatment Hardness, HV30*

Mn As received microstructure

WQ†

350 oC for 30 min, WQ

597±2

605±5

467±7

Ni1 As received microstructure

WQ


493±5

647±8

426±4

Ni2 As received microstructure

WQ


536±6

669±7

423±9

*Hardness values are averaged over 10 tests

†WQ water quench

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Table 5 Tensile* and Fracture Toughness† Properties

Alloy YS

MPa

UTS

MPa

Elongation

%

RA

%

Kmax

MPa m

Jmax

MPa m

KJmax

MPa m

Mn 1167 1790 13 44

Ni1 1150 1725 14 55 125 0.114 160

Ni2 1100 1625 14 59 128 0.134 174

*YS yield strength; UTS ultimate tensile strength; RA reduction of area.

†Kmax stress intensity factor at maximum load; Jmax J-integral at maximum load;

KJmax stress intensity values calculated from Jmax values.

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Table 6 Charpy Impact Test Results*

Alloy Test temperature, oC Impact energy, J

Mn 20 34±1

-40 31±1

Ni1 20 58±2

-40 46±1

-120 34±2

Ni2 20 50±3

-40 43±1

-120 25±3

*Each value is a mean of six tests

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Figure 1 Manufacturing route

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a Mn alloy as received microstructure; b Ni1 alloy as received microstructure;

c Ni1 alloy as tempering at 200 oC for 2 hrs. microstructure; d Ni2 alloy as

received microstructure

Figure 2 Scanning electron micrographs of microstructures in the designed

alloys

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a Ni1 alloy; b Ni2 alloy

3 Typical bright field images of microstructures formed by bainitic ferrite and

films of retained austenite

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a Mn alloy, room temperature; b Ni1 alloy, room temperature; c Ni2 alloy, room

temperature

4 Fracture surfaces of Charpy impact specimens tested at different

temperatures. Scanning electron micrographs

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5 Properties of mixed microstructures of bainitic ferrite and austenite, versus

those of quenched and tempered (QT) low-alloy martensitic alloys and

maraging steels

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a Mn alloy; b Ni1 alloy; c Ni2 alloy

6 Plot of hardness values as a function of tempering temperature

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a Discrete carbide particles precipitate inside a martensite plate in Mn alloy; b

None sings of recovery occur in the bainitic microstructure of Ni1 alloy; c Bright

field and d dark field images reveal that retained austenite is still intact in Ni2

alloy

7 Transmission electron micrographs of microstructures obtained by

tempering at 400 oC for an hour

design of high-strength bainitic...

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